Spinel compounds as multivalent battery cathodes: a
systematic evaluation based on ab initio calculations
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Citation
Liu, Miao, Ziqin Rong, Rahul Malik, Pieremanuele Canepa,
Anubhav Jain, Gerbrand Ceder, and Kristin A. Persson. “Spinel
Compounds as Multivalent Battery Cathodes: a Systematic
Evaluation Based on Ab Initio Calculations.” Energy Environ. Sci.
8, no. 3 (2015): 964–974. © 2015 Royal Society of Chemistry
As Published
http://dx.doi.org/10.1039/C4EE03389B
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Cite this: Energy Environ. Sci., 2015, 8,
964
Spinel compounds as multivalent battery cathodes:
a systematic evaluation based on ab initio
calculations†
Miao Liu,‡a Ziqin Rong,‡b Rahul Malik,b Pieremanuele Canepa,b Anubhav Jain,a
Gerbrand Cederb and Kristin A. Persson*a
Batteries that shuttle multivalent ions such as Mg2+ and Ca2+ ions are promising candidates for achieving
higher energy density than available with current Li-ion technology. Finding electrode materials that
reversibly store and release these multivalent cations is considered a major challenge for enabling such
multivalent battery technology. In this paper, we use recent advances in high-throughput first-principles
calculations to systematically evaluate the performance of compounds with the spinel structure as
multivalent intercalation cathode materials, spanning a matrix of five different intercalating ions and
seven transition metal redox active cations. We estimate the insertion voltage, capacity, thermodynamic
stability of charged and discharged states, as well as the intercalating ion mobility and use these
properties to evaluate promising directions. Our calculations indicate that the Mn2O4 spinel phase based
on Mg and Ca are feasible cathode materials. In general, we find that multivalent cathodes exhibit lower
voltages compared to Li cathodes; the voltages of Ca spinels are 0.2 V higher than those of Mg
compounds (versus their corresponding metals), and the voltages of Mg compounds are 1.4 V higher
than Zn compounds; consequently, Ca and Mg spinels exhibit the highest energy densities amongst all
Received 24th October 2014
Accepted 16th December 2014
the multivalent cation species. The activation barrier for the Al3+ ion migration in the Mn2O4 spinel is
very high (1400 meV for Al3+ in the dilute limit); thus, the use of an Al based Mn spinel intercalation
DOI: 10.1039/c4ee03389b
cathode is unlikely. Amongst the choice of transition metals, Mn-based spinel structures rank highest
www.rsc.org/ees
when balancing all the considered properties.
Broader context
The high price and limited volumetric capacity of the lithium ion battery (LIB) challenges its application in electric vehicles and portable electronics. Multivalent
batteries, such as those utilizing Mg2+ or Ca2+ as the working ions, are promising candidates for beyond LIB technology due to the increase in volumetric capacity
and reduced cost. In the present work, we use rst-principles calculations to systematically evaluate the theoretical performance of the spinel structure host with
the general formula AB2O4 across a matrix of chemical compositions spanning A ¼ {Al, Y, Mg, Ca, Zn} and B ¼ {Ti, V, Cr, Mn, Fe, Co, Ni} for multivalent battery
applications. The evaluation incorporates screening on voltage, capacity, thermodynamic structural and thermal stability as well as ion mobility and discusses
the results in the context of available host structure sites, preference of the intercalating cation, and the oxidation state of the redox-active cation. Overall, the
Mn2O4 spinel phases paired with Mg2+ or Ca2+ emerge as the most promising multivalent cathode materials. As the rst comprehensive screening of multivalent
intercalation compounds across size, valence, and redox-states of the involved cations, our work is intended to provide guidance for future theoretical as well as
experimental multivalent cathode development and design.
Introduction
To support the rapidly growing energy storage demands of
future technologies such as electric vehicles, improved battery
a
Environmental Energy Technologies Division, Lawrence Berkeley National Laboratory,
CA 94720, USA. E-mail: kapersson@lbl.gov
b
The Department of Materials Science and Engineering, Massachusetts Institute of
Technology, Cambridge, MA 02139, USA
† Electronic supplementary
10.1039/c4ee03389b
information
(ESI)
‡ M. Liu and Z. Rong contributed equally to this work.
964 | Energy Environ. Sci., 2015, 8, 964–974
available.
See
DOI:
technologies are needed. Lithium (Li) ion batteries with good
energy density, rechargeability and cycle life1 have been used to
advance portable consumer electronics and recently, electric
vehicles. However, further advancement of Li ion technology
faces limits on the energy density of electrode materials,2,3
safety, and high cost. A multivalent battery technology, where
an intercalation cathode host is paired with a metal anode, has
the potential to store energy at signicantly lower cost and
volume.4–6 For example, Mg metal has much higher volumetric
capacity (3833 mA h cm3) than graphite (800 mA h cm3)7,8 or
even lithium metal (2046 mA h cm3). Furthermore, as many
known intercalation hosts are limited by the available cation
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insertion sites rather than by redox capability, one can expect
improvements using multivalent intercalation on the cathode
side. For instance, Mg2+ carries two charges per ion while
maintaining an ionic radius as small as that of a Li+ ion, hence
the charge storage capability is doubled at the same cation
volumetric concentration. In addition, the natural abundance
of multivalent elements, such as Mg and Ca, is signicantly
higher than that of Li (the atomic abundance of Mg is 104
times larger than Li in the earth crust), which may decrease the
cost of materials and guarantee supply even for multi-fold
increases in the size of the energy storage market.
To date, there have been limited examples demonstrating
the feasibility of rechargeable multivalent batteries, and among
them, most of the focus has been on Mg technology. In the
seminal work of Aurbach et al., it was demonstrated that Mg2+
can intercalate into the Chevrel phase Mo6S8 at a potential of
1.0–1.3 V and a maximum charge capacity of 135 mA h g1 for
several hundreds of cycles.6,9 Other materials have shown initial
promise for multi-valent intercalation, such as layered V2O5
where Mg2+ is inserted into the V2O5 inner layer spacing.10,11
Ca2+ has also been shown to intercalate into V2O5, supplying a
voltage of 3.0 V and an initial capacity of 450 mA h g1.12
Other materials claimed to accommodate Mg2+ insertion
include MoO3,13–15 TiS2 in both the cubic phase and layered
phases,16 NbS3,17 graphitic uoride,18 and CoSiO4.19 However,
due to limited Mg mobility in the host and possibly concurrent
water and/or proton insertion, the cycling stability of these
materials has so far been insufficient, resulting in rapid
decomposition of the host material.20,21
From the limited experimental studies performed to date,
the feasibility of a battery technology based on multivalent
intercalation is not yet clear. As the cathode will be a critical
component of such a technology, it is important to assess the
feasibility of multivalent insertion cathodes. In this work, we
present a systematic computational study of multivalent intercalation within a xed spinel-based host structure, using a set of
seven redox-active cations and spanning size and valence
differences between a set of intercalating {Ca, Zn, Mg, Al and Y}
cations to establish design trends and guidelines for future
experimental work.
m) structure
The spinel (prototype MgAl2O4, space group Fd3
provides an excellent candidate for this study, encompassing a
family of materials with the general formula AB2O4. The A and B
ions are tetrahedrally and octahedrally coordinated by oxygen,
respectively (Fig. 1). The B octahedrons form a network with
percolating empty sites interconnecting in three directions.
Spinel LiMn2O4 was rst prepared by Thackeray et al.22 and
exhibits excellent performance as a cathode for Li intercalation
with a voltage of 3–4 V versus Li metal.23 The properties of the
spinel structure are tunable, for example it has been observed
that partially replacing Mn with Ni increases the voltage to 4.7
V,24,25 and mixing Mn with Co and/or Cr increases the voltage
even higher to 5 V.25–27
Experiment shows that spinel LiMn2O4 can be electrochemically converted to MgMn2O4 in aqueous Mg(NO3)2 electrolyte, and exhibits Mg2+ reversible intercalation/
deintercalation.28 Spinel ZnMnO2 has also demonstrated Zn2+
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Energy & Environmental Science
Fig. 1 The spinel crystal structure where the ‘A’ atoms occupy the
tetrahedral sites, and the ‘B’ atoms occupy the octahedral site.
Throughout this paper, the ‘A’ atoms are multivalent intercalating ions
selected from the set {Mg2+, Ca2+, Zn2+, Y3+, Al3+}, and the ‘B’ atoms
are transition redox-active ions, selected from the set {Ti, V, Cr, Mn, Fe,
Co, Ni}.
insertion/deinsertion, providing a capacity of 210 mA h g1 for
50 cycles.29 Recently, it was shown that Mg2+ can intercalate/
deintercalate into spinel-type Mn2O4 with a retained capacity of
155.6 mA h g1 aer 300 cycles in 1 mol dm3 MgCl2 aqueous
electrolyte,30 suggesting that further development could lead to
a viable cathode with good energy density. Against this background it is intriguing to broadly consider a multivalent spinel
cathode with possible intercalating cations A ¼ {Mg, Ca, Zn, Al,
Y} that could theoretically produce a higher capacity than its Li
counterpart due to the greater charge carried by each ion.
Hence, in this paper, our aim is to evaluate the cathode
performance of spinel phases for multivalent intercalation; we
computationally evaluate the feasibility of a matrix of spinel
compounds with different redox ion species and intercalating
cation species. The redox ion was selected from the set {Ti, V,
Cr, Mn, Fe, Co, Ni}, and the intercalating cation was selected
from the set {Mg, Ca, Zn, Al, Y}. By substituting the cations in
the A and B sites, respectively, we created 35 charged/discharged topotactic pairs and performed rst-principles density
functional theory (DFT) calculations for each of these pairs. We
evaluate cathode performance through such quantities as the
capacity, average voltage, energy density, and intercalating
cation mobility. We also evaluate the thermodynamic structural
and thermal stability. The detailed methodology can be found
in previous literature.31
Methods
We use the Vienna ab initio soware package (VASP)32 to
perform the density functional theory calculations, with the
projector augmented-wave method33 to describe the ion–electron interactions and the generalized gradient approximation
(GGA)34 within the Perdew–Burke–Ernzerhof (PBE) framework35
as the exchange–correlation functional. The calculation
parameters are the same as those adopted by the Materials
Project36 and as implemented in the pymatgen soware
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package,37 which have been previously tested to be appropriate
to study Li-intercalation cathode materials.31 In the calculations, the U–J parameters to correct for non-cancellation of the
self-interaction error in the d orbitals of the redox active species
are set to UV ¼ 3.25 eV, UCr ¼ 3.7 eV, UMn ¼ 3.9 eV, UFe ¼ 5.3 eV,
UCo ¼ 3.32 eV, and UNi ¼ 6.45 eV.38 The primitive spinel unit cell
as illustrated in Fig. 1 is used for voltage and stability calculations. Brillouin zone sampling is performed on a 5 5 5 grid
in k-space. Throughout this work, the cell shape, volume and
atomic positions are relaxed, unless otherwise stated. All
magnetic ions are initialized ferromagnetically.
All calculations in this paper assume that the transition
metal host framework ‘B2O4’ remains structurally invariant
during the operation of the battery (i.e., during intercalation
and de-intercalation of ‘A’ cations). Additionally, we assume
that the host can be synthesized with little/no disorder and
remains that way during the operation of the cell. We
acknowledge that spinels are well known to show varying
degrees of cation disorder that may impact important material
properties relevant for battery operation such as the activation
energies and voltages reported herein. However, in the interest
of providing a preliminary view of what is possible in multivalent systems, we have simplied our calculations and analysis.
The voltages of the compounds can be obtained from the
difference in the total energy between the charged and discharged phases following Aydinol et al.39,40 The average voltage
can be calculated as V ¼ DE/nz, where DE ¼ (Echarge + EMV Edischarge) denotes the total energy change in the reaction, Echarge
and Edischarge are the energy of the charged and discharged
compounds respectively; EMV is the energy of multivalent
intercalating species in metal form; n is number of intercalating
atoms participating in the reaction; and the z represents the
oxidation state of the intercalant. We adopt the units of eV and e
for DE and nz, respectively, so that no normalization factor (i.e.,
Faraday's constant) needs to be introduced into the equation.
We estimate the thermodynamic stability of the phases by the
energy above the convex hull of stable phases, which is the
energy released by decomposing the compound to the most
stable combination of compounds at the same overall composition.41,42 The energy above the hull is always a non-negative
number with the unit of eV per atom. The detailed procedure for
the computation of the energy above the hull can be found in
previous literature.41,42 As explained later in this paper, the
thermal stability was determined by evaluating the critical
chemical potential at which O2 gas becomes favorable according to the methodology presented in Ong et al.43 Since the
entropy of a reaction is dominated by the gas entropy and the
entropy of the solid phase at room temperature,38 for the any
reactions containing molecular O2, we use the corrected O2
chemical potential to include the well-known O2 DFT calculation error38 as well as the PDV contribution to the oxygen
enthalpy43 by comparing with the experimental thermodynamic
data44 for O2 at 0.1 MPa at 298 K throughout the paper. For the
thermal stability calculation, the temperature effect has been
taken into account by adjusting the entropy term (TDS) of the
O2 chemical potential at given temperature.45 A more detailed
description can be found in the ESI.†
966 | Energy Environ. Sci., 2015, 8, 964–974
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The calculations were automatically executed and analyzed
using the FireWorks soware package.46 In this work, hundreds
of DFT calculations are performed across 70 compounds to
generate the thermodynamic stability and thermal stability
data.41,43
Activation barriers were calculated with the nudged elastic
band (NEB) method47 using the GGA-PBE functional.34,35 A U
term was not included in these calculations as NEB is difficult to
converge with GGA+U due to pronounced metastability of
electronic states along the ion migration path. Furthermore,
while GGA+U clearly improves the accuracy of redox reactions,31
there is no conclusive evidence that GGA+U performs better in
predicting cation migration.48–52 The minimum energy paths
(MEP) in the NEB procedure were initialized by linear interpolation of 8 images between the two fully relaxed end-point
geometries, and each image is converged to <1 104 eV per
super cell. The MEPs were obtained in both the high vacancy
limit and dilute vacancy limit, i.e. one mobile species per unit
cell or one vacancy per unit cell. To ensure that ctitious
interactions between the diffusing species are removed, a 2 2
2 supercell of the primitive cell was used, for which the interimage distance is never less than 8 Å.
Results and discussion
The average intercalation voltage was calculated from the
reaction energy B2O4 + A / AB2O4 for the matrix of intercalating A ¼ {Mg, Ca, Zn, Y, Al} ions and redox active transition B ¼
{Ti, V, Cr, Mn, Fe, Co, Ni} metal cations. Fig. 2 shows the
resulting voltage vs. the gravimetric capacity, where the color
and shape of a data point indicate the intercalating ion, and the
redox-active transition metal is given next to each data point.
Voltages are referenced to the bulk metal of the intercalating
ion, e.g., Mg metal for MgMn2O4 and Zn metal for ZnMn2O4.
As expected from the electrochemical series, and evidenced
in Fig. 2, multivalent compounds have lower voltages than Li
Fig. 2 The computed average voltage vs. gravimetric capacity for
intercalation of A ¼ Zn, Ca, Mg, Y and Al in various M2O4 spinels up to
composition AM2O4. The redox-active metal is marked next to each
point. Dashed curves show the specific energy of 600 W h kg1, 800 W
h kg1 and 1000 W h kg1, respectively. The spinel LiMn2O4 and olivine
LiFePO4 data points are also marked on the plot for comparison.22,36,43
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cathodes. The Li spinel usually exhibits a voltage between 4 and
5 V,22,26,53 whereas we nd that for multivalent spinel cathodes,
the voltage is always less than 4 V versus the corresponding
metal. Most Mg and Ca intercalation voltages vary between 2
and 4 V, which is lower than the values for Li (e.g., calculated
voltage of LiMn2O4 is 0.7 V higher than CaMn2O4);22,54
however, considering the additional charge carried by multivalent cations, a multivalent spinel cathode can still exhibit a
signicantly higher energy density than the corresponding Li
version. For example, the gravimetric capacity of LiMn2O4 is 143
mA h g1,22,36 whereas the gravimetric capacity of MgMn2O4 is
almost double 270 mA h g1, which more than makes up for
the slightly lower voltage. Hence, of the considered intercalating
ions, Al, Y, Ca and Mg are all viable candidates from the
perspective of energy density. However, the voltage of the Zn
spinel compounds ranges between 1.3 V and 2.5 V, which even
in the best case scenario amounts to approximately 600 W h
kg1, about equal to the specic energy of LiFePO4.36,43
The voltage for each multivalent intercalant is plotted as a
function of the active redox metal in Fig. 3(a): the bi-valent ions
Fig. 3 (a) The calculated voltage of each spinel phase for the reaction
B2O4 + A / AB2O4 as a function of the redox-active transition metal
and intercalating cation. The different colors denote different intercalating species as specified by the legend. The black triangle point
indicates the data corresponding to the spinel LiMn2O4;22,36 (b) and (c):
for Y and Al the voltage for separate 3+/2.5+ and 4+/3+ redox reactions is compared to the overall voltage for the 4+/2.5+ redox change
corresponding to the full intercalation range, as in (a).
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Ca, Mg, and Zn follow a common trend as the redox couple is
varied, different from the voltage trend of the tri-valent intercalants Al and Y. The difference originates largely from the
different valence state of the transition metal in the discharged
state. Insertion of the bi-valent cations induces a change in
redox state from 4+ to 3+, whereas the tri-valent cation corresponds to a redox change from 4+ to 2.5+ for insertion into
AB2O4.
In general, the Ca spinel has the highest voltage, followed by
the Mg spinel compounds, Y compounds, Al compounds, and
Zn compounds, in that order. For all the redox active cations {B
¼ Ti, V, Cr, Mn and Ni} the voltage of the Mg compounds is
lower than that of the Ca compounds by 0.2 V, and the voltages of Mg compounds are 1.4 V higher than Zn compounds.
The three bi-valent intercalants show the same trend of voltage
verses redox active metal: Ti2O4 always has the lowest voltage
among the transition metals considered. V2O4 is the second
lowest one but 1.2 V higher than Ti2O4. Mn2O4 is 0.3 V
higher than V2O4, Co2O4 is 0.6 V higher than Mn2O4, and
Ni2O4 is slightly higher than Co2O4 by 0.1 V. The Cr2O4 and
Fe2O4 spinels have the highest voltage, respectively 0.6 V and
0.9 V higher than Mn2O4. Bhattacharya et al. found a similar
trend for the Li insertion voltage in spinels.54 We nd that Li
insertion54 occurs on average at about 0.7 V higher voltage
than Ca insertion and 0.9 V higher than Mg insertion.
Comparing this with the aqueous electrochemical series (E0Li ¼
3.04, E0Ca ¼ 2.86, E0Mg ¼ 2.37, E0Zn ¼ 0.76) we nd that the
voltages are ordered according to the electrochemical series of
the intercalating metal ion. However, while the voltage shi
between Li and Mg is close to what is expected, the voltage
reduction in moving from Li to Ca in the solid state is considerably larger than expected from the electrochemical series.
This is likely due to the fact that the intercalant enters a tetrahedral site in the spinel, which for Ca is not nearly as favorable
as for Mg and Li, and thus reduces the Ca intercalation voltage
from what one would expect from the electrochemical series.
Experimentally, Li has indeed been found to exhibit a higher
voltage than Mg. In the Chevrel phase Mo6S8, the voltage
difference between Li and Mg insertion is 1.0–1.2 V.9,55,56 For
V2O5, the Li voltage is usually 0.2 V higher than that of Mg.57–59
Because the data in Fig. 3(a) for trivalent cations averages the
voltage over both the 3+/2.5+ and 4+/3+ redox couples, the
intermediate 3+ states of the transition metals were calculated
for the trivalent intercalants to investigate the impact of
different redox states. Fig. 3(b) and (c) show the calculated
voltage of the different redox pairs for the Al and Y spinel
compounds, respectively. As expected, the 3+/2.5+ redox pairs
exhibit a lower voltage compared to the 4+/3+ reactions, in good
agreement with existing literature.2 In particular, when
restricting focus to the 4+/3+ redox reaction, the Al and Y spinel
compounds follow largely the same trend as Ca, Mg and Zn as
shown in Fig. 3(a).
The capacity of cathode materials is important as it strongly
inuences the overall energy density of a cell. Fig. 4(a) shows the
volumetric capacity of each AB2O4 spinel as a function of the
redox-active species and intercalating cation. The volumetric
capacity of all the cathodes is higher than that of a Li spinel
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Fig. 4 The calculated (a) volumetric multivalent capacity and (b)
volume change of the spinel structure as a function of the redox-active
cation, assuming intercalation to composition AB2O4.
cathode at the same cation concentration due to the extra
charge carried by each of the multivalent intercalants. Not
surprisingly, Al3+ leads to the highest capacity density, while
Ca2+ has the lowest. For a xed valence of the intercalant, the
volumetric capacity follows the ionic size of the intercalating
cation. Hence, the volumetric capacities of Al3+ compounds are
higher than Y3+ compounds by approximately 300 A h L1. For
bi-valent cations, the capacities of Mg and Zn compounds are
almost the same, consistent with the similar ionic size of Mg2+
or Zn2+.
Fig. 4(b) presents the volume expansion associated with the
intercalation of each multivalent ion as function of the redox
metal. Volume change is an important parameter as it needs to
be accommodated at the particle, electrode and cell level. At the
particle level, large volume changes can lead to particle fracture
and loss of contact. The total volume change associated with
intercalation is the combined result of the intercalant insertion
and the transition metal reduction, and can be very small for
some Li-insertion systems.60,61 The contribution from the
intercalating ion will depend on its size and charge. For
example, Y3+ leads to larger volume increase than Al3+ as the ion
is much larger. The effect of charge cannot as easily be extracted
from Fig. 4(b) as the +3 cations also cause a larger reduction of
the transition metal than the 2+ cations. Reduction of the
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transition metals generally leads to an increase in volume,
although the magnitude of the change depends on the nature of
the metal-d orbital that is being lled. Filling of t2g orbitals,
which are to rst order non-bonding,40 tends to cause only a
small increase in volume, while the anti-bonding eg orbitals
lead to a larger volume change.60 This explains why in general
the early transition metals, such as Ti and V exhibit lower
volume changes. They have several unoccupied t2g orbitals,
which are available for reduction. On the other hand, Mn and
Ni-based spinels show the largest volume changes as their
reduction occurs by lling one or more eg orbitals. For both
Mn3+ and Ni3+ this volume increase is compounded by the fact
that these ions are Jahn–Teller active which, in its anharmonic
form, leads to additional volume increase.62 The combined
small size and high charge of Al3+ lead to almost zero volume
change for several spinels.
For Mg, Zn and Al insertion, the magnitude of the volume
change normalized by the capacity is very similar to the volume
changes observed for Li insertion compounds, and hence is not
likely to lead to any practical design problems. For Ca and Y the
volume change is larger – up to 30% increase in some cases.
The thermodynamic stabilities of the charged and discharged state are important considerations for possible cathode
materials, as they may inuence the cycle life as well as the
synthesizability of the compounds. Thermodynamic stability
can be measured by the driving force for a compound to separate into its most stable combination of compounds. From rst
principles, this is determined by comparing the energy of the
compound with the convex energy hull of all ground states in
the relevant phase diagram (see the ESI† for a detailed explanation). Fig. 5(a) and (b) provide this energy above the hull of
each compound. The ground state hulls were determined from
all the calculated compounds in the Materials Project database.36 A smaller energy above the hull implies that the material
has a greater chance of being stable,63 e.g. at synthesis and upon
cycling.
The Mg and Zn spinel phases (except ZnTi2O4) are all quite
stable, exhibiting an energy above hull less than 0.011 eV per
Fig. 5 (a) The calculated thermodynamic stability of the AB2O4 spinel
compounds as a function of the intercalating ion (vertical axis) and
redox metal (horizontal axis). (b) The energy above hull of the charged
state which is calculated as the formation energy difference between a
compound and the convex hull (see ESI† for a detailed explanation).
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atom. Such a small value of the decomposition energy falls well
within the accuracy of our calculations63 or within relative
changes between competing phases due to nite temperature
effects. The Ca spinel structures are more thermodynamically
unstable compared to the equivalent Mg and Zn compounds.
This is consistent with Ca2+ normally preferring coordination
environments larger than tetrahedral, though it is not excluded
that this coordination can be achieved through electrochemical
intercalation. The Y spinel compounds are the most unstable
phases among all the candidates, due to the large ionic size of
Y3+. The Al3+ spinels are also relatively unstable in the discharged state. Among the chemistries of the charged host
spinel, Mn2O4 forms the most stable structure. The high energy
above hull for Fe2O4 suggests the difficulty to synthesize such a
phase. Indeed, Fe4+ is only known to exist in ternary and higher
component compounds64 where the formation energy is lowered by the interaction with other cations. In terms of phase
stability across both charged and discharged states, we
conclude that MgMn2O4, CaMn2O4 and ZnMn2O4 provide the
best opportunities.
A way to gauge the intrinsic safety of a potential cathode
material is by the thermal stability of the compound against O2
release. The thermal stability can be estimated by calculating
the temperature at which O2 gas release is predicted thermodynamically (see ref. 43 and the ESI† for more details on the
methodology). Fig. 6 shows the calculated O2 amount released
as a function of temperature for the charged spinel compounds
(e.g. for B2O4), determined by the equilibrium chemical potentials at which O2 release can be expected. At low temperature the
oxygen release is likely limited by kinetics and the decomposition temperatures should only be used to rank compounds
relative to their oxidation strength. Clearly, the data in Fig. 6
indicates that Fe2O4 and Ni2O4 are highly oxidizing, and are
unlikely to be stable in their stoichiometric conguration at
room temperature. Co2O4 decomposes to 2/3(Co3O4 + O2) at a
temperature slightly above 100 C. The Cr2O4 and Mn2O4
spinels are predicted to decompose to Cr2O3 + 1/2O2 and Mn2O3
+ 1/2O2 at 285 C and 342 C, respectively. Since the Mn spinel
Fig. 6 The calculated thermodynamic O2 evolution diagram of
charged host spinel compounds as a function of temperature. The
vertical axis denotes how much O2 is predicted to be released from the
compound per formula unit of B2O4 as the material decomposes.
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Energy & Environmental Science
exists as a metastable phase in some Li batteries, the Cr and Mn
spinel phases should operate well at room temperature. Finally,
the V2O4 and Ti2O4 spinels are predicted as fairly stable against
O2 release as indicated by their higher decomposition temperatures; 876 C and 1646 C, respectively. In summary, Ti2O4,
V2O4, Cr2O4 and Mn2O4 are expected to exhibit superior thermal
stability against O2 release among the considered spinel
compounds. Comparing the calculated voltage (Fig. 3) and
thermodynamic stability (Fig. 5), a stable discharged phase and
unstable charged phase naturally lead to higher voltage, and
vice versa (as expected). For example, fairly unstable Fe2O4 and
Ni2O4 result in a higher voltage, while Mn2O4 generally results
in lower voltages.
The previous voltage, capacity and stability results for
multivalent cathode materials show great potential to go
beyond current Li-ion. However, a major remaining challenge is
overcoming the sluggish diffusion expected for multivalentions. The slower diffusion of high-valent cations has been
attributed to the stronger cation–anion interaction, which
makes migrating 2+ or 3+ ions more difficult than moving 1+
ions, though no quantitative information is available on mulitvalent ion diffusion.5 Hence, we calculate the migration energy
barriers of the various multivalent ions (A ¼ Mg2+, Zn2+, Ca2+,
Al3+, and also Li+ for comparison) in the spinel structure AB2O4
(B ¼ Mn, Co, Ni, Cr) from rst-principles as shown in Fig. 7. The
Nudged Elastic Band method was used in both high vacancy
limit and dilute vacancy limit corresponding a single migrating
intercalant or a vacancy in an empty host or fully intercalated
structure, showing the upper and lower limit of the intercalant
migration activation barrier. To provide a general idea of how
the migration energy barrier affects the diffusivity, a rough
estimation can be given as follows: a migration barrier of 525
meV corresponds to a typical ionic diffusivity of 1012 cm2 s1
at room temperature, which approximately represents the lower
limit for reasonable charge and discharge time (2 h in micronsize active particles). Also, a 60 meV increase (decrease) in the
migration energy corresponds to an order of magnitude
decrease (increase) in the intercalant migration activation
barrier. With smaller particle sizes a somewhat larger barrier
can be tolerated: every order of magnitude in size reduction
allows for two orders of magnitude smaller diffusion constant.
Hence for 100 nm particulates, migration barriers up to z645
meV may be acceptable.
In spinel Mn2O4, which has been successfully commercialized for use in Li-ion batteries, Al3+ displays the highest diffusion barrier of 1400 meV. Among the divalent cations, Zn2+
(850–1000 meV) and Mg2+ (600–800 meV) have the highest
barriers, while Ca2+ is comparable to Li+ (400–600 meV). The
migration barriers obtained for Li+ in M2O4 (M ¼ Mn, Co, Ni, Cr)
all lie within 400 to 600 meV in the empty lattice limit, in good
agreement with rst-principles Li mobility calculations performed by Bhattacharya et al. in the spinel Li1+xTi2O4 system
(1 < x < 1).65 Excluding the CrO2 spinels and some of Ca-containing spinels, the migration barrier at high vacancy limit is
always higher compared to the dilute vacancy limit, agreeing
with Bhattacharya et al. that migration barrier is reduced by
nearly 300 meV as Li is intercalated from the Li-decient to the
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Fig. 7 Computed minimum energy paths for migration of different intercalants between the tetrahedral sites in the spinel B2O4 (B ¼ Mn, Co, Ni,
Cr) at the high vacancy limit (solid line) and dilute vacancy limit (dotted line), i.e. one mobile specie per supercell (2 2 1 of primitive cell).
in Mn2O4, we could not converge the NEB for Al3+ due to the very
large forces along the transition path, which is usually symptomatic of a very high barrier. The divalent barriers vary
signicantly with the chemical nature of the intercalant: Zn2+
(800–1000 meV), Mg2+ (600–800 meV), and Ca2+ (200–500
meV). Although Ca2+ migration appears to be facile in our
calculations, only in Mn2O4 does Ca2+ prefer to occupy the
tetrahedral site as opposed to the octahedral site (which can be
observed in Fig. 7 as the energy along the migration path
becomes negative when Ca is near the octahedral site in the Ni,
Cr, and Co spinel). For this case, the migration barrier should
be measured as the energy increase from the octahedral site to
the maximum along the path (see Fig. 7). We nd that for all
fully intercalated phases, the tetrahedral sites are more stable
than the octahedral sites, which indicates a cross-over in site
preference with concentration for some of the intercalating
cations.
Qualitative summary of multivalent spinel compounds based
on multiple performance metrics, such as voltage, specific energy,
thermodynamic stability of charged and discharged phases, thermal
stability and intercalant mobility. The favorable (unfavorable) properties are represented with light (dark) color.
Fig. 8
Li-rich limit (from 600 meV for Li migration in Ti2O4 to 300
meV for vacancy migration in LiTi2O4).65 From kinetic Monte
Carlo simulations, the room temperature self-diffusivity of Li
was shown to span 1010 to 109 cm2 s1 between Li0.5Ti2O4
and LiTi2O4, in good agreement with the excellent experimental
rate-capability typically observed in Li spinel cathodes. Except
970 | Energy Environ. Sci., 2015, 8, 964–974
Discussion
Technical extrapolations of projected multivalent chemistries to
the cell level have shown that multivalent intercalation is one of
the few technologies that can outperform Li-ion batteries in
terms of energy density. In this paper we have used rst principles calculations, well established in Li-cathode
research,2,3,43,63,66 to evaluate the properties of multivalent-ion
intercalation in spinel structures with different chemistries. We
evaluated average insertion voltage, stability in the charged and
discharged state, volume change upon intercalation, oxidation
strength of the charged cathode, and the mobility of the multi-
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Paper
valent cations. Such rst principles screening is important for
this new eld as multi-valent ion electrochemistry is not well
established and, due to incompatibility between electrolytes
and electrode materials, it can be difficult to obtain unambiguous experimental results on the performance of a specic
cathode material.4,67,68 We found that the insertion voltage of
Ca2+, Mg2+, Zn2+, Al3+ and Y3+ against their respective metal
anodes in general follows the electrochemical series but with
quantitative variations due the nature of the site preference of
the intercalating ion. For example, Ca insertion voltages in
spinels are lower than expected, as Ca in general prefers higher
coordination than tetrahedral (commonly the most favorable
site in the spinel structures). The Fe2O4 and Ni2O4 spinels are
unlikely to function across the full capacity range due to the
highly oxidizing and unstable nature of their fully charged
states. As expected, Ti2O4 spinels have low insertion voltage for
most intercalants, and in addition, are fairly unstable. The V2O4
and Cr2O4 spinels are also fairly unstable in the charged state,
and the V spinel exhibits a low insertion voltage for all intercalants besides Mg and Ca. From the perspective of stability, the
most promising spinel chemistry is Mn2O4 as it is stable in the
charged state and fairly stable in the discharged state for several
intercalating ions. Intercalant mobilities are generally low due
to the high activation energies when transitioning between the
tetrahedral and octahedral site, though they are clearly not only
controlled by charge. For example, among the divalent ions Zn
mobility is inferior to Mg, and Ca may have fairly good mobility
in the spinel. Al3+ intercalation into the spinel structure can
likely be excluded from consideration even though it has the
peculiar feature that volume changes upon insertion are very
small. The activation barrier for motion is very high for Al3+ in
Mn2O4 and its discharged spinels are all highly unstable.
Somewhat surprisingly, and in contrast to experimental
claims,29 we nd that mobility of Zn in the spinel structure is
very low, which in addition to its low insertion voltage, should
exclude this system from further consideration for high energy
density cathodes. The large intercalating ions such as Y3+ and
Ca2+ are interesting, though the Y-spinels become unstable in
the discharged limit. Both of these ions have reasonable
insertion voltages and Ca2+ has better than expected migration
barriers, due to its relative instability in the tetrahedral site.
While this effect lowers the voltage from what would be expected by considering the electrochemical scale, it also seems to
lower the migration barrier for motion. This may be similar to
the more general principle that high energy defects in materials
oen have higher mobility.
Considering all computed properties, Mn2O4 spinels are
particularly interesting due to their stability (Fig. 8). Among the
divalent cations, both Mg2+ and Ca2+ may potentially be mobile
in the spinel structure, warranting further experimental and
computational investigation (particularly at small particle
sizes). Mixed spinel structures may provide a further promising
avenue, as the Ni4+/3+ and Co4+/3+ show higher voltage than the
Mn4+/3+ redox couple and compounds such as LiNiMnO4 are
known to combine higher operating voltage with a relatively
stable charged state.24,25 Improving the diffusivity should be a
focus for all multivalent cathode materials.
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Energy & Environmental Science
Finally, a concern with spinels, as with all intercalation
materials, is the occurrence of cation disorder. In the extreme
case, normal spinels can convert to inverse spinels with part of
the transition metals on the tetrahedral site. While in more
common layered materials, lowering of the intercalant mobility
by cation disorder is well understood through the contraction of
the interlayer slab space by disorder,69 we are not aware of an
equivalent study for spinels. Given the 3D and more rigid nature
of the framework and the 3D diffusion network, one would
expect spinels to be more tolerant to cation disorder. Nonetheless, we have performed a preliminary investigation into the
driving force inverse spinel formation by adopting the same
methodology as in Bhattacharya et al.54 Based on small supercell
calculations, we nd that the normal spinel is always energetically favorable for Ca, Mg, Zn compounds (See Fig. S2 in the
ESI†). Further information can be obtained from the study of
Burdett et al. who elucidated that cation disorder is strongly
correlated to the relative size of the A and B ions.70 Hence, Li
and Zn prefer to form normal spinels, and we do not expect
cation disorder. Ca is large and has the possibility to occupy
octahedral sites that forms inverse spinel. Similarly, Mg may
show cation site disorder for Ni, Fe, Co when synthesized at
high temperature. However, in the case that well ordered
spinels cannot be formed at high temperature synthesis
conditions, it may be still be possible to create the normal
ordered spinel phase by chemical or electrochemical delithiation of the lithium spinels28,30 and inserting multivalent cations.
In summary, we have performed systematic calculations to
screen for and discover improved multivalent cathode materials
using the spinel structure as a general host. On the basis of all
property calculations, the spinel Mn2O4 is found to be a superior candidate with Ca2+ and possibly Mg2+ as mobile cations. It
is our hope that our work provides a general guide and standard
for future theoretical as well as experimental multivalent
cathode development and design.
Acknowledgements
This work was intellectually led and fully supported by of the
Joint Center for Energy Storage Research (JCESR), an Energy
Innovation Hub funded by the U. S. Department of Energy,
Office of Science, Basic Energy Sciences. Work at the Lawrence
Berkeley National Laboratory was supported by the Assistant
Secretary for Energy Efficiency and Renewable Energy, under
Contract no. DEAC02-05CH11231. We also thank the National
Energy Research Scientic Computing Center (NERSC) for
providing computing resources. The Materials Project (BES
DOE Grant no. EDCBEE) is acknowledged for infrastructure and
algorithmic support.
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